Ionic to Electronic Transport in Ba3Ti3O6(BO3)2 under Reducing

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Ionic to Electronic Transport in BaTiO(BO) under Reducing Atmosphere Jean-Marie Doux, Khang Hoang, Olivier Joubert, Jonathan Hamon, Florian Massuyeau, and Eric Quarez ACS Appl. Energy Mater., Just Accepted Manuscript • DOI: 10.1021/acsaem.7b00124 • Publication Date (Web): 31 Jan 2018 Downloaded from http://pubs.acs.org on February 3, 2018

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ACS Applied Energy Materials

Ionic to Electronic Transport in Ba3Ti3O6(BO3)2 under Reducing Atmosphere Jean-Marie Doux,† Khang Hoang,‡,* Olivier Joubert,† Jonathan Hamon,† Florian Massuyeau,† Eric Quarez†,* †

Institut des Matériaux Jean Rouxel (IMN), Université de Nantes, CNRS, 2 rue de la Houssinière, B.P. 32229, 44322 Nantes Cedex 3, France ‡

Department of Physics, North Dakota State University, Fargo, North Dakota 58108, United States

ABSTRACT: The influence of a reducing atmosphere on transport properties of undoped and doped oxyborate Ba3Ti3O6(BO3)2 has been studied using various experimental methods and first-principles calculations. We find that the electrical transport mechanism in the material changes from being ionic under oxidizing and soft reducing conditions to mainly electronic under extreme reducing conditions; consequently, the total conductivity is increased by about 200 times. The change is caused by the emergence of electron polarons, associated with the reduction of Ti 4+ to Ti3+, as the predominant highly mobile current-carrying defects. Hybrid-density-functional electronic structure and defect calculations provide a novel theoretical framework for understanding the observed transport properties, including the ionic-toelectronic transition and the effects of doping. We also find that Ba3Ti3O6(BO3)2 is stable in the extreme reducing atmosphere and the Ti4+/Ti3+ reduction and re-oxidation reactions are reversible. This mixed ionic-electronic conductor can thus have applications in solid-oxide fuel cell anodes or anode composites.

Keywords: oxyborate; mixed ionic-electronic conductor; defect physics; SOFC; anode 1. INTRODUCTION Solid-oxide fuel cell (SOFC) is one of the promising technologies to respond to the increasing energy demand while potentially being one of the most environmentally friendly.1,2 It has the advantages of providing a high energy-conversion efficiency and long-term stability as well as the possibility to operate with different kinds of fuels. Good electrochemical performances can be achieved with conventional SOFC anodes based on NiYSZ.3 However, performance degradation of Ni-based anodes due to redox cycling instability and Ni agglomeration has been reported.4,5 Besides, when used with hydrocarbon fuels, the performance also gets reduced because of carbon deposition and sulphur poisoning.6–8 Numerous studies have been carried out to conceive Ni-free anode materials. Complex oxides such as LaxSr1–xTiO3– and related materials are some of the promising candidates.9–11 Indeed, in these titanates, the Ti4+ ion can be reduced to Ti3+ in reducing atmosphere,12 providing a means for electronic transport. Searches for new anode materials should thus also include transitionmetal-containing compounds. Recently, a new class of complex materials–oxyborates– was highlighted and considered for SOFCs.13,14 Among these materials, Ba3Ti3O6(BO3)2, see Fig. 1, undoped or doped with subvalent or supervalent elements, was stud-

ied in detail. It was found that this Ti-containing material is an oxygen-ion conductor in atmosphere ranging from oxidizing (with the oxygen partial pressure pO2 = 1 atm) to soft reducing (pO2 ~ 10–5 atm) conditions. The ionic transport and effects of doping were understood on the

Figure 1. Structural representation of Ba3Ti3O6(BO3)2. Large (green) spheres are Ba, medium (orange) Ti, small (gray) B, and smaller (red) O. The three inequivalent O sites, O1 (connecting Ti octahedra along the c-axis), O2, and O3 (linked to B), are indicated.

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basis of the material's defect physics and chemistry obtained from first-principles calculations.14 It remains to be explored, however, if the material exhibits an electronic contribution in addition to the ionic conduction when subjected to an extreme reducing atmosphere (pO2 ~ 10–20– 10–23 atm). In this case, it may be used at the anode side of SOFCs. On the fundamental side, a theoretical framework for understanding possible changes in the electrical transport mechanism due to a change in the environment, e.g., via temperature and the oxygen partial pressure, is highly needed for the study of SOFC materials. Here, we report a combined experimental and computational study of transport properties in undoped and doped Ba3Ti3O6(BO3)2 under extreme reducing conditions (pO2 ~ 10–20 atm). The electrical transport mechanism in the material is found to change from being purely ionic to mainly electronic in going from oxidizing and soft reducing conditions to extreme reducing conditions. Experiments are also carried out to investigate the stability, reversibility, and kinetics of the Ti4+/Ti3+ reduction and re-oxidation reactions, to determine the conditions for preparing fully reduced samples, to estimate the Ti3+/Ti4+ proportion in the fully reduced samples, to study the effects of doping on the amount of Ti3+, and to determine the position of the Fermi level. The observed transport properties are analyzed and discussed on the basis of a theoretical framework developed based on results from hybrid-density-functional electronic structure and defect calculations. Finally, potential applications of Ba3Ti3O6(BO3)2 -based materials in SOFCs are also discussed. 2. EXPERIMENTAL AND COMPUTATIONAL SECTION 2.1. Powder syntheses. Powders of Ba3Ti3O6(BO3)2 and Ba3Ti3–xMxO6(BO3)2 (M = Mg2+, In3+, Nb5+, Ta5+; x = 0.03, 0.06) were prepared by solid-state reaction in air according to the experimental conditions described in Ref. 14. Briefly, high purity oxides and carbonates (> 99.8%) were mixed in appropriate molar ratio and annealed at 950°C for 24 h. 10 mol.% of H3BO3 was added in excess to compensate for boron evaporation during heat treatment. The heating and cooling rate of the furnace was 120°C h–1. The obtained powders were then manually ground and another heat treatment step was carried out to increase powder purity. These samples will be called hereafter unreduced samples. 2.2. X-ray characterization. X-ray powder diffraction (XRD) patterns were recorded as a function of temperature in air or in N2/H2 5% using a Bruker D8 Advance equipped with an Anton Paar HTK 1200N hightemperature attachment. Data were collected in the Bragg–Brentano geometry with a Cu anode X-ray source and CuK radiation filtered by means of a Ni foil (CuK1: = 1.54060 Å, CuK2: = 1.54439 Å). The detector was a 1-D position-sensitive detector (Vantec detector). The heating rate was chosen to be 0.1 °C s–1. The experiments were performed in the 10–80° 2 range with a step of 0.016 ° and an acquisition time of 1.2 s per step. Longer XRD patterns were recorded at room temperature (RT) in air using a Bruker D8 Advance diffractometer working in the Bragg–Brentano geometry with a Cu anode X-ray

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source, a focusing Ge (111) primary monochromator (CuK1:  = 1.54060 Å) and a 1-D silicon-strip PSD (LynxEye detector). The conditions used were 10–110° 2 range with a step of 0.012° and an acquisition time of 4.92 s per step. The cell parameters refinements were carried out using FULLPROF15 in the full pattern matching mode with WinPLOTR.16 2.3. Pellet preparation. Dense ceramics of ~9.4 mm diameter and ~1 mm thickness were obtained by uniaxial compression at 150 MPa of the powders followed by heat treatment at 950°C for 24 h in air or in Ar/H2 5%. All pellets were then polished with fine abrasives. For all sintered pellets, a density higher than 92% of the theoretical density was obtained.14 Gold ink was used as the electrical contact and was painted on both sides of the pellets that were calcined in vacuum at 800 °C for 10 h (unreduced pellets) or in Ar/H2 5% (reduced pellets). 2.4. UV-Visible (UV-Vis) Spectroscopy. Diffuse reflectance spectra were recorded in the range of 250–800 nm with a 1 nm step at RT using a Lambda 1050 UV-visible spectrometer from Perkin-Elmer equipped with a 150 mm integrating sphere. The reflectance data were converted to the Kubelka-Munk function F(R) using the equation, F(R) = [(1 – R)2/2R], where F(R) is proportional to the absorption coefficient and R is the reflectance at a given wavelength. The band-gap value was obtained by plotting (F(R) × h)1/2 as a function of photon energy E and determined from the intersection of the energy axis and the tangent to the linear part of the curve. About 500 mg of powder was used for each measurement. 2.5. Seebeck Coefficient Measurements. Seebeck coefficients were obtained using a home-made apparatus consisting of both cold and hot blocks. The hot block is a cylindrical piece of brass with a heating resistor connected to a DC voltage generator. The cold block is made of copper and was chosen large enough to stay at RT in the working temperature interval. Sintered pellet is placed between both blocks and a K-type differential thermocouple allows the measurement of the temperature difference between each side of the pellets. Electrical contacts, to measure the thermoelectric voltage, are placed on each block. To ensure a good electrical and thermal contacts, silver paste was spread between blocks and each side of the pellet, and the whole assembly was pressed. Pellets were heated with a maximal temperature difference of about 10 K between each side of the pellet, then heating was switched off and measurements were carried out during cooling. The thermoelectric voltage and temperature difference were measured with a Keithley 182 and an Agilent 3458A voltmeters controlled with LabVIEW. 2.6. Electrochemical Characterization. Electrochemical impedance spectroscopy was performed with a frequency response analyzer Solartron 1260.14 Each spectrum was recorded at UDC = 0 V, with a signal amplitude of 100 mV from 4 MHz to 0.1 Hz frequency range. Data were analyzed using the ZView 2 software.17 The conductivity measurements were performed during cooling, from 750 to 300 °C every 25 °C after a stabilization time of 20 min. Two different atmospheres were employed: an oxidizing


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atmosphere supplied by compressed air from a network (pO2 ~ 0.2 atm) and a reducing atmosphere obtained by a cylinder containing Ar/H2 5% (pO2 ~ 10–20 atm). In both cases, gas was dried out by passing through a column filled up with silica gel. In the following, log (σ) = f(1000/T) is plotted instead of log (σT) = f(1000/T) to facilitate the reading of the σ conductivity values. However, all values of activation energies Ea were calculated based on the log (σT) = f(1000/T) plots.

all calculations, the plane-wave basis-set cutoff was set to 500 eV and spin polarization was included.

2.7. Thermogravimetric (TG) Analysis. TG analyses were performed with a Netzsch STA 449F3 Jupiter. Approximately 200 mg of finely grinded powder was heated from 25 to 900 °C and cooled down to RT at a rate of 1 K min–1, in Ar/H2 5% or O2. The Netzsch Proteus software was used to extract the data.

where Etot(Xq) and Etot(bulk) are the total energies of a supercell containing the defect and that of the perfect material. µi is the atomic chemical potential of species i that has been added (ni > 0) or removed (ni < 0) to form the defect. The chemical potentials of Ba, Ti, B, and O can be chosen to reflect the experimental conditions.14 µe is the electron chemical potential, i.e., the Fermi level, referenced to the VBM in the bulk (EV). In Eq. (1), µe can be regarded as a variable; however, the actual Fermi level is determined by the charge neutrality condition that involves all defects and any other charge carriers present in the material.26

2.8. X-ray Photoelectron Spectroscopy (XPS). Spectra were collected using a Kratos Axis Ultra and Nova X-ray photoelectron spectrometers using monochromatic Al Kα X-rays (1486.6 eV). The base pressure in the analysis chamber was 10–8 Pa, and the analyzed area was 700 μm × 300 μm. Survey scans were recorded with a step size of 0.5 eV and a pass energy of 80 eV, and high resolution spectra of Ti 2p, Nb 3p and Nb 3d core levels were recorded with a step size of 0.1 eV and a pass energy of 40 eV. For valence-band spectra, the pass energy was 20 eV. Surface charging from photoelectron emission was neutralized using a built-in system. Spectra were calibrated using the C 1s peak at 284.8 eV. The Ag 3d5/2 FWHM is 0.47 eV giving an instrumental resolution about 0.1 eV. Data analysis of the core levels (Ti 2p, Nb 3d) was conducted with the standard CasaXPS software. Each 2p (Ti) and 3d (Nb) core-level spectrum is constituted of spinorbit doublets (2p3/2 2p1/2 and 3d5/2 3d3/2). The photoemission contributions were fitted using the Tougaard function for the background. The Ti 2p1/2 peak for each species is constrained to a fixed energy above the Ti 2p3/2 peak. The intensity ratio of the Ti 2p3/2 and Ti 2p1/2 peaks are also constrained to 2:1. The position of the valence-band maximum (VBM) was determined using the "Edge Down" function in CasaXPS software. 2.9. First-Principles Calculations. The calculations were based on density-functional theory (DFT), using the Heyd-Scuseria-Ernzerhof (HSE06) screened hybrid functional,18,19 the projector augmented wave method,20,21 and a plane-wave basis set, as implemented in the Vienna Ab Initio Simulation Package (VASP).22–24 We used the standard values of the mixing parameter (25%) and the screening length (10 Å). The calculations for bulk Ba3Ti3O6(BO3)2 were carried out using a 3×3×7 k-point mesh. We modelled intrinsic point defects and impurities (extrinsic defects) using hexagonal 1×1×3 (60-atom) supercells. Integrations over the supercell Brillouin zone was carried out using a Γ-centered 2×2×2 k-point mesh. Structural relaxations were performed with the HSE06 functional and the force threshold was chosen to be 0.02 eV/Å. Additional defect calculations were carried out using larger, 2×2×3 (240-atom), supercells and the DFT+U method25 with the effective U value of 4.0 eV applied on the Ti 3d states. In

A defect can be characterized by its formation energy. Defects with lower formation energies will form more easily and occur with higher concentrations. The formation energy of a defect X in charge state q is defined as26

E f ( X q )  Etot(X q )  E tot (bulk)   ni μi  q( E V  μ e ),

(1)

i

3. RESULTS AND DISCUSSION 3.1. Unreduced, Partially Reduced, and Fully Reduced Samples. In order to study undoped and doped Ba3Ti3O6(BO3)2 samples in reducing atmosphere, three kinds of samples were prepared and characterized, hereafter throughout the article referred to as unreduced, partially reduced, and fully reduced samples. Unreduced samples are those that were prepared according to the procedures described in Sec. 2.1 and at no time submitted to a reducing atmosphere; all steps from powder synthesis to pellet annealing were carried out in air (i.e., oxidizing atmosphere). Partially reduced samples were obtained from annealing unreduced samples at 730°C for 5 h in Ar/H2 5% inside the furnace before running conductivity measurements (The reason for the choice of these experimental conditions will be made clear in Sec. 3.5). Fully reduced samples were obtained from unreduced samples by performing back and forth steps of powder annealing in Ar/H2 5% and reflectance measurements of the resulted powders until no evolution of the reflectance was observed. Figure 2 illustrates how Ba3Ti3O6(BO3)2 and Ba3Ti3– 5+ xMxO6(BO3)2 (x = 0.06; M = Nb ) change, from being unreduced to fully reduced, when the unreduced samples were subjected repeatedly to steps of grinding and annealing at 850°C in Ar/H2 5%. After each annealing cycle, the reflectance of the resulting powders was measured by UV-Vis spectroscopy; see Fig. 2a. Before reduction, the reflectance is about 90% between 400 and 800 nm; it then decreases progressively as the powders are subjected to further annealing in the reducing atmosphere. After several annealing cycles, the reflectance reaches a minimum at about 20%. As the reflectance does not evolve after two successive UV-Vis measurements, we assume that the sample attains its maximum reduction state and refer to it as being fully reduced.


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Figure 2. (a) Reflectance as a function of wavelength for Ba3Ti3O6(BO3)2 and Ba3Ti3–xMxO6(BO3)2. The number of hours spent in the reducing atmosphere is indicated in the case of the undoped sample. (b) Color of the pellets at RT before reduction and after full reduction steps. (c) Transformed Kubelka-Munk function vs. energy and obtained optical band gap (Eg).

It is noted that the color of the material changes from white grey for the unreduced samples to dark blue for the fully reduced samples (Fig. 2b), suggesting that a certain amount of Ti4+ ions has been reduced to Ti3+. This behavior is similar to that previously observed in Ti-based materials such as TiO227,28 and BaTiO3.29,30 The indirect optical band gap31,32 of the undoped and Nb-doped samples was determined to be Eg = 3.33 eV (Fig. 2c), which is in good agreement with that obtained in HSE06 calculations (3.47 eV).14 As reduction time increases, the band-gap value of the undoped sample decreases progressively until reaching a minimum value at 3.05 eV. For the Nb-doped sample, the band gap at the end of the reduction process is 2.96 eV, slightly smaller than that in the case of the undoped sample. 3.2. n-Type Conductivity of Undoped and Doped Samples in Ar/H2 5%. Figure 3 shows the conductivity of partially reduced Ba3Ti3O6(BO3)2 and Ba3Ti3–xMxO6(BO3)2 (M = Mg2+, In3+, Nb5+, Ta5+; x = 0.03, 0.06) samples measured in Ar/H2 5%. The conductivity of unreduced samples of the same compositions in air, previously shown to be purely ionic,14 is also included for comparison. Clearly, the conductivity level of the partially reduced samples in Ar/H2 5% is higher than that of the unreduced ones in air by about 200 times whereas the activation energies are divided by a factor of 2. Here, given the change in color of the reduced samples (Sec. 3.1) and the dramatic conductivity increase, it is strongly suggested that Ti4+ → Ti3+ reduction occurs upon subjecting the undoped and doped samples to extreme reducing atmosphere. We note that the reduction was not observed for Ba3Ti3O6(BO3)2 under soft reducing conditions.14 Also, the dramatic change in the conductivity (and, presumably, the Ti4+ → Ti3+ reduction) was not observed in other Ti4+-containing compounds such as BaIn0.8Ti0.2O2.6 and BaIn0.6Ti0.2Yb0.2O2.6 even when subjected to the Ar/H2 5% atmosphere.33

ties are, respectively, lower and higher than that of the undoped one; see Fig. 3. The Nb- and Ta-doped samples show the best conductivity level and are thus the most promising for applications. The effects of doping on the conductivity will be discussed in Sec. 3.3. We note that further characterizations are performed only on the Inand Nb-doped samples, used as representative cases of subvalent and supervalent doping, respectively.

Figure 3. Conductivity of partially reduced Ba3Ti3O6(BO3)2 and Ba3Ti3–xMxO6(BO3)2 measured in Ar/H2 5% (closed symbols) and corresponding activation energy Ea. The conductivity of unreduced samples of the same compositions in air (open symbols)14 is included for comparison.

The conductivity levels in Ar/H2 5% of the samples doped with subvalent (Mg, In) and supervalent (Ta, Nb) impuri-


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Figure 5. Total and projected electronic density of states (DOS) of Ba3Ti3O6(BO3)2. The zero of the energy (at 0 eV) is set to the highest occupied state. Figure 4. Thermoelectric voltage vs. temperature difference for Ba3Ti3O6(BO3)2 and Ba3Ti3–xMxO6(BO3)2. The sign and the value of the Seebeck coefficient are given by S = – dV/dT.

Seebeck coefficient measurements were also carried out for the fully reduced Ba3Ti3O6(BO3)2 and Ba3Ti3– 3+ 5+ xMxO6(BO3)2 (x = 0.06, M = In , Nb ) samples to determine the type of charge carriers. The Seebeck coefficients at RT are all negative, as seen in Fig. 4, indicating that carriers in the fully reduced samples are negatively charged and the material exhibits an n-type conductivity (see also Secs. 3.3 and 3.8 for an in-depth discussion of the origin and further evidence of the n-type conductivity in the material). Note that attempts to determine Seebeck coefficients of unreduced samples using our homemade apparatus were unsuccessful. 3.3. Fundamental Understanding from FirstPrinciples Calculations. In order to understand the origin of the observed n-type conductivity and effects of doping, we investigated the electronic structure and the formation and migration of relevant defects in Ba3Ti3O6(BO3)2. Figure 5 shows the electronic density of states of the compound. We focus primarily on the electronic structure near the band edges as it is relevant to the material's transport and electrochemical properties.34 It is evident that the valence-band top is predominantly composed of the O 2p states, whereas the conductionband bottom is predominantly the empty Ti 3d states with the conduction-band minimum (CBM) formed by the Ti t2g states. Titanium is stable as Ti4+ with a zero magnetic moment. The electronic structure indicates that the reduction of Ba3Ti3O6(BO3)2 would involve adding electrons to the Ti 3d orbitals (thus turning Ti4+ into Ti3+) and the oxidation would involve removing electrons from the O 2p orbitals (though, in this work, we are more interested in re-oxidation of the reduced Ba3Ti3O6(BO3)2 which involves turning Ti3+ back into Ti4+). This has important consequences on defect formation and hence defectmediated transport as discussed in detail below.

Figure 6a shows the formation energies of oxygen hole polarons (ηO+) and electron polarons at the titanium site (η–). ηO+ occurs through the removal of an electron from the supercell that results in a localized electron hole at one of the O1 lattice sites (see Fig. 1) or, equivalently, an O– ion with a calculated magnetic moment of 0.71 μB. The local lattice environment is distorted in the presence of the localized hole; the Ti–O bonds associated with ηO+ have a bond length of 2.05 Å (compared to 1.93 Å in the bulk material). The energy of ηO+ is lower than that of a free hole by 0.63 eV. η–, on the other hand, occurs through the addition of an electron to the supercell that leads to the formation of a localized electron at one of the Ti4+ sites and turns the Ti4+ ion into Ti3+ with a calculated magnetic moment of 0.79 μB. The local environment of the localized electron is also distorted with an average Ti3+–O bond length of 2.02 Å, larger than that of the Ti4+– O bonds in the bulk (1.97 Å). The energy of η– is lower than that of a free electron by 0.12 eV. The formation of the polarons can be understood in terms of the electronic structure presented above. To facilitate our discussions, the results for other relevant intrinsic point defects are also reproduced in Fig. 6a. These defects are oxygen vacancies VO (at the O1 site), oxygen interstitials Oi, and barium vacancies VBa. We note that, among different possible charge states, only the following configurations are the true charge states of the mentioned defects: VO2+, Oi2–, and VBa2–. Other (nominal) charge states are, in fact, complexes consisting of these elementary configurations and hole or electron polaron(s). For example, the neutral VO, nominally denoted as VO0, is a complex of VO2+ and two η–; similarly, VBa0 is a complex of VBa2– and two ηO+. It is also noted that the formation energies presented in Fig. 6a are obtained with the atomic chemical potentials chosen to reflect the experimental conditions under which unreduced Ba3Ti3O6(BO3)2 was prepared, as detailed in Ref. 14. Under these conditions, the Fermi level of the material is at μeint = 2.15 eV where the charge neutrality is maintained.14 This position is determined predominantly by the dominant charged defects VO2+ and VBa2–. The polarons and those


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Figure 6. (a) Formation energies of relevant intrinsic defects occurred during the synthesis of Ba3Ti3O6(BO3)2 in air, plotted as a function of Fermi level from the VBM to the conduction-band minimum (CBM); μeint, marked by the vertical dotted line, is the position of the Fermi level determined by the charge neutrality condition that involves all intrinsic defects under the synthesis conditions. The slope indicates charge state q: positively (negatively) charged defects have positive (negative) slopes. (b) Formation energies of relevant intrinsic defects occurred during the reduction process; μeint,red is the Fermi-level position determined by the charge neutrality condition under the reducing conditions. (c) Formation energies of In and Nb impurities, where the blue (green) arrow indicates the direction in which the Fermi level would be shifted when the concentration of Nb (In) at the Ti site is higher than that of the lowest-energy positively (negatively) charged intrinsic point defect.

defect complexes containing polarons have much higher energies than VO2+ and VBa2–; see Fig. 6a. As discussed in Ref. 14, VO2+, VBa2–, and (to a lesser extent) Oi2– are expected to occur with high concentrations in Ba3Ti3O6(BO3)2. The above mentioned defects are expected to remain in the material after synthesis and act as preexisting athermal defects in subsequent experiments.14 When subjecting the material to an extreme reducing atmosphere, such as that under which the reduced Ba3Ti3O6(BO3)2 samples were prepared, more defects will be activated. Under the new experimental conditions, which can be translated into a new set of atomic chemical potentials, the defect landscape will be different from that discussed earlier. Figure 6b shows the formation energies of the intrinsic defects that occur during the reduction process. Here, we assume that the system exchanges only oxygen species (and electrons) with the environment and there is no Ba and/or B loss. The formation energies of the O-related defects were obtained with μO = –2.99 eV, the oxygen chemical potential at 730°C and 10–20 atm35–the conditions under which the partially Ba3Ti3O6(BO3)2 samples were prepared (see Sec. 3.1). We find that the formation energy of VO2+ (Oi2–) is now lower (higher) by 1.52 eV than that obtained under the oxidizing conditions (Fig. 6a), whereas that of the polarons stays constant as it is independent of the atomic chemical potentials. The position of the Fermi level is now μeint,red = 2.98 eV, determined predominantly by VO2+ and η–. With this new Fermi-level position, the formation energy of η– is now much lower, making it easy to form and occur in a high concentration. η– and VO2+ can occur in the form of the neutral complex VO0

with a formation energy of only about 1 eV, see Fig. 6b. Note that the formation of VO0 in the material is equivalent to having oxygen removed from the material. Given the presence of hydrogen in the reducing environment, we also investigated hydrogen as an impurity in Ba3Ti3O6(BO3)2 and found it is most stable as the positively charged interstitial Hi+ bound to O at the O1 site with an O–H bond length of 0.97 Å. However, the hydrogen species is expected to be released with the removed oxygen at high temperatures, as evidenced by the mass loss occurring during the reduction process discussed later in Sec. 3.8. Thus far, we have shown that η– (i.e., Ti3+) can form in Ba3Ti3O6(BO3)2 and that it can occur, together with VO2+, with a high concentration under an extreme reducing atmosphere. η– and VO2+ remain in the material after the reduction process and will act as preexisting athermal defects in subsequent measurements carried out on the reduced samples. Additional η– and VO2+ will also be thermally activated during the conductivity measurements in a reducing environment such as Ar/H2 5%. Let us now discuss how these defects contribute to electronic transport in the material. It is noted that polarons migrate in the material through hopping. In Ba3Ti3O6(BO3)2, the migration of η– from one Ti site to another can be described by the transfer of its lattice distortion. Following the method described in, e.g., Ref. 36 and references therein, the energy barrier (Em) of η– was estimated to be 0.10 eV along the Ti–O1–Ti path (c-axis). This value is slightly lower than those estimated in the DFT+U calculations using 2×2×3 supercells, Em = 0.17 eV (along the Ti– O1–Ti path, i.e., c-axis) and 0.24 eV (along the Ti–O2–Ti


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path, i.e., b-axis). Note that the migration barrier of VO2+ was reported to be higher than 1 eV,14 i.e., the oxygen vacancy can be regarded as immobile compared to the polaron. Given that difference between VO2+ and η– and the fact that their formation energies at μeint,red are almost equal under the extreme reducing atmosphere (Fig. 6b), the ionic contribution to the total conductivity must be much smaller than the electronic one. Furthermore, that ionic part is expected to be comparable to the ionic conductivity of unreduced Ba3Ti3O6(BO3)2 in air14 as we find that the formation energy of VO2+ at μeint in air (Fig. 6a) is comparable to that at μeint,red in Ar/H2 5% (Fig. 6b). Our results, therefore, indicate that the conductivity observed in the reduced Ba3Ti3O6(BO3)2 (Fig. 3) is mainly electronic, in contrast to the purely ionic nature of the conductivity of the unreduced samples measured in air. The electronic conduction occurs through hopping of η– and is thus n-type as observed in the Seebeck measurements (Fig. 4). The estimated migration barrier of η– is smaller than the measured activation energies Ea, as expected. As discussed in detail in Refs. 36,37, the effective activation energy is dependent on the ratio between the athermal and thermally activated defects. The lower limit Ea = Em corresponds to the situation when the currentcarrying defects are predominantly athermal (e.g., preexisting in the samples), typically occurring at low temperatures. The upper limit Ea = Ef + Em often occurs at high temperatures where the current-carrying defects are predominantly thermally activated.36,37 We now turn our attention to the effects of doping on the conductivity measured in a reducing atmosphere. Figure 6c shows the calculated formation energies of NbTi and InTi impurities in Ba3Ti3O6(BO3)2. In the present work, the energies were obtained by setting μNb = –7.00 eV and μIn = –1.00 eV. This choice of the atomic chemical potentials is somewhat arbitrary; however, it does not affect our discussions and conclusions as we are interested only in the stable charge states and not in the solubility of the dopants. It is evident from the figure that Nb and In are stable as NbTi+ (i.e., Nb5+ at the Ti4+ lattice site) and InTi– (i.e., In3+ at the Ti4+ site) at μeint–the position of the Fermilevel of the undoped material under the synthesis conditions. These defect configurations are also the true charge states of NbTi and InTi. For example, the neutral NbTi, nominally denoted as NbTi0, is in fact a complex of NbTi+ and η–; InTi0 is a complex of InTi– and ηO+, whereas InTi+ is a complex of InTi– and two ηO+. Our calculations thus confirm the oxidation state of Nb (+5) and In (+3) in Ba3Ti3O6(BO3)2 assumed in Ref. 14 and earlier in the current work. We note that the concentration of the dopants in doped Ba3Ti3O6(BO3)2 samples is 1–2% and thus should be higher than that of the lowest-energy defects in the unreduced material. In this case, the Fermi level in Nbdoped Ba3Ti3O6(BO3)2 is no longer determined by VO2+ and VBa2– (i.e., μeint in Fig. 6a) but by NbTi+ and VBa2– and is shifted toward the CBM (the direction of the blue arrow in Fig. 6c). As a result of the Fermi-level shift, the formation energy of η– will get reduced. The Fermi level in

In-doped Ba3Ti3O6(BO3)2, on the other hand, is shifted toward the VBM (the green arrow) where the charge neutrality condition is maintained predominantly by VO2+ and InTi–. This Fermi-level shift will increase the formation energy of η–. Experimental evidence for the Fermi-level shift will be presented in Sec. 3.8. Note that somewhat similar argumentation related to shifting the Fermi level has also been discussed in the context of mass and charge transport in hydrogen storage38,39 and battery40 materials. The effects of supervalent and subvalent doping on the conductivity, as presented in Fig. 3, can now be explained as the following: Since the formation energy of η– in the Nb-doped sample is lower than that in the undoped one, the creation of η– upon exposing the material to the reducing atmosphere is more effective, particularly at the beginning of the reduction process. As a result, the reduced Nb-doped sample has a higher concentration of athermal preexisting η– defects than in the undoped one during the conductivity measurements. This explains the lower activation energy and higher conductivity observed at low temperatures in Fig. 3. Experimental evidence of the enhanced Ti3+ concentration will be presented in Secs. 3.6–3.8. In the case of the In-doped sample, similar argumentation can be made. Here, due to the higher formation energy of η– for the reduction process to begin with, the reduction at the beginning of the process is less effective. As a result, the concentration of preexisting η– defects in the reduced In-doped sample is lower than in the undoped sample, effectively increasing the activation energy in the In-doped sample as seen in Fig. 3. Finally, Mg and Ta are expected to be stable as MgTi2– and TaTi+ and would behave similar to InTi– and NbTi+, respectively. The conductivity of the Mg-doped sample is even lower than that of the In-doped one, see Fig. 3. This is due to the fact that the Mg impurity is doubly negative, which makes the reduction process even less effective than in the case of In. The difference between the conductivity levels of TaTi+ and NbTi+ is much smaller, see Fig. 3, as expected because they have the same effective charge. 3.4. Stability and Reversibility of the Reduction/ReOxidation Reaction. Since Ba3Ti3O6(BO3)2 decomposes into BaTiO3 and Ba4Ti12O27 at 975°C in air,14 thermal XRD experiments whose purpose is to study the stability and reversibility of the Ti4+/Ti3+ reaction were carried out up to 900°C. Figure 7 shows the XRD patterns from RT to 900 °C and the final XRD pattern at RT of the unreduced Ba3Ti3O6(BO3)2 sample in N2/H2 5% and the fully reduced Ba3Ti3O6(BO3)2 sample in air. We find that the samples are stable up to 900°C, both in N2/H2 5% and in air, and the final XRD pattern shows that the Ti4+/Ti3+ reduction reaction is reversible without degradation of the sample. We also note that the white gray powder at the beginning of the XRD experiment in N2/H2 5% is notably darkened at the end of the experiment, consistent with the occurrence of the Ti4+ → Ti3+ reduction reaction; whereas the dark blue powder at the beginning of the experiment in air is notably lighter at the end of the experiment, indicating the occurrence of the Ti3+ → Ti4+ re-oxidation reaction.


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Figure 7. XRD patterns from RT to 900 °C and the final XRD pattern at RT of (a) unreduced Ba3Ti3O6(BO3)2 in N2/H2 5% and (b) fully reduced Ba3Ti3O6(BO3)2 in air. Inset: Lattice parameter change as a function of temperature.

Refinements in the 𝑃6̅2𝑚 space group based on 12 h XRD patterns at RT in air give cell parameters a = 8.7121(1) Å, c = 3.9302(1) Å, and V = 258.337(5) Å3 for unreduced Ba3Ti3O6(BO3)2,14 and a = 8.71117(10) Å, c = 3.93222(5) Å, and V = 258.417(5) Å3 for fully reduced Ba3Ti3O6(BO3)2. The cell parameters of the unreduced and fully reduced Ba3Ti3O6(BO3)2 samples are thus almost identical, indicating that the reduction process has negligible influence on the overall structure. In this case, it is likely that effects of the oxygen loss on the cell parameters are compensated by those of the reduction in which a certain amount of Ti4+ (rVI = 0.605 Å) is replaced by the larger Ti3+ ions (rVI = 0.670 Å), especially when the amount of Ti3+/oxygen loss is relatively small (which is, indeed, the case as will be discussed in Secs. 3.7 and 3.8). The thermal expansion coefficients, of great interest to determine if different SOFC components are compatible,13,41,42 are shown in the insets of Fig. 7. Between 25 and 900 °C, they are found to be similar to that previously obtained in the case of unreduced samples annealed in air (~ 12×10–6 K–1)14 and very close to that of the stabilized zirconia electrolyte reference material (~ 11×10–6 K–1).43 3.5. Kinetics of the Reduction and Re-Oxidation Reactions. In order to study the kinetics of the Ti4+/Ti3+

redox reactions, conductivity measurements were carried out as a function of time for partially reduced samples of Ba3Ti3O6(BO3)2 and Ba3Ti3–xMxO6(BO3)2 (x = 0.06; M = In3+, Nb5+) at constant temperature (635°C). Figure 8a shows the conductivity upon several step changes of the oxygen partial pressure pO2. It can be seen from the figure that the conductivity change when one switches from Ar/H2 5% to air is more abrupt than from air to Ar/H2 5%, indicating that the re-oxidation process is faster than the reduction one. It is noted that a duration longer than 25 h in Ar/H2 5% is needed to recover the initial conductivity level. This behavior may be explained as the following: When subjected to the atmosphere change from Ar/H2 5% to air, the material at/near the surface of the pellet is quickly oxidized even when the interior is still reduced, leading to an impediment to electronic transport (and hence a sharp drop in the conductivity) due to a lack of η– (i.e., Ti3+) for hopping. When atmosphere is switched from air to Ar/H2 5%, the reduction occurs first at/near the surface. Since the reduction process must reach the bulk to have efficient pathways for η– hopping, this process is expected to be slow. Further discussion of the difference between the reduction and re-oxidation processes is presented in Sec. 3.7.


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Figure 8. (a) Conductivity vs. time upon step changes of pO2 for partially reduced Ba3Ti3O6(BO3)2 and Ba3Ti3–xMxO6(BO3)2. (b) Normalized conductivity vs. time upon a step change of pO2 from air to Ar/H2 5%. Inset: t1, time at which 99% of the normalized conductivity is reached, at different temperatures; the dashed line is to guide the eyes.

Figure 8b shows the normalized conductivity upon one step change from air to Ar/H2 5% at 700°C. The evolution of the conductivity was recorded as a function of time (t), where t = 0 corresponds to the time at which the atmosphere was switched from air to Ar/H2 5%. The conductivity was found to be almost constant after 5 h for all three samples. We, therefore, carried out heat treatment for 5 h in Ar/H2 5% at a temperature higher than 700°C (about 730°C) in the furnace to ensure that samples were close to their maximum reduction state before running the conductivity measurements. This is the reason for our choice of 5 h and 730°C as mentioned in Sec. 3.1–the experimental conditions under which partially reduced samples were prepared. The inset of Fig. 8b shows the values of t1, time at which 99% of the normalized conductivity is reached, at different temperatures. In the studied temperature range, t1 follows a linear trend and decreases as temperature increases, i.e., the reduction process is faster at higher temperatures, as expected. In the next section (Sec. 3.6), we will discuss in more detail the influence of the reduction state on the electrical conductivity and activation energy. 3.6. Dependence of the Conductivity on the Reduction State. During the operation of a SOFC, the anode is subjected to extremely low oxygen partial pressures (pO2 ~ 10–20–10–23 atm). The Ba3Ti3O6(BO3)2 material used as an anode would thus be reduced. Besides, one typically reduces SOFC anodes before running the cells.44,45 As discussed in Sec. 3.1, it can take a long time for the material to be fully reduced, however. It is, therefore, important to know how the conductivity level depends on the reduction state. Figure 9 shows the conductivity of the partially reduced vs. fully reduced Ba3Ti3O6(BO3)2 and Ba3Ti3–xMxO6(BO3)2 (x = 0.06; M = In3+, Nb5+) in Ar/H2 5%. Interestingly, we find that although the conductivity levels of the samples are quite distinct at low temperatures they are almost the same at high temperatures. The effec-

tive activation energy is in the range 0.41–0.62 eV for the partially reduced samples and 0.29–0.36 eV for the fully reduced samples. The results thus indicate that it is not really necessary to have Ba3Ti3O6(BO3)2 fully reduced beforehand for SOFC anodes operating at a typical operating temperature (700°C). The above results can be understood based on the theoretical framework presented in Sec. 3.3, particularly regarding the dependence of the effective activation energy

Figure 9. Conductivity of partially reduced (closed symbols) and fully reduced (open symbols) Ba3Ti3O6(BO3)2 and Ba3Ti3– xMxO6(BO3)2 in Ar/H2 5%. Ea is the activation energy.


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on the ratio between athermal and thermally activated defects. The activation energy is lower in the fully reduced and/or Nb-doped samples, as expected because of the higher concentration of athermal η– defects. Because of the dependence on the concentration of preexisting defects, the activation energy is sensitive to how the samples were prepared. It is noted that the measured activation energy of the fully reduced Nb-doped sample is about 0.29 eV, comparable to the calculated migration barrier of η–. In this case, the current-carrying defect η– is likely to be predominantly athermal and hence Ea ~ Em(η– ).36,37 The situation is different at high temperatures where thermally activated η– defects are expected to be predominant. In this case, the activation energy is insensitive to the methods used in samples preparation. Indeed, the conductivity levels for all samples reported in Fig. 9 become comparable at about 700°C or higher.

the defects that occur during the reduction process (see Fig. 6b) are included.

3.7. Mass Loss (Gain) from the Reduction (ReOxidation) Reaction. Figure 10 shows TG analyses of unreduced or reduced Ba3Ti3O6(BO3)2 and Ba3Ti3– 5+ xMxO6(BO3)2 (x = 0.06; M = Nb ) samples. Clearly, there is mass loss during the heating of the unreduced samples in Ar/H2 5%; see the blue curves. This phenomenon can be understood in terms of the defect landscapes discussed in Sec. 3.3 according to which oxygen vacancies occur in Ba3Ti3O6(BO3)2 (equivalently, oxygen release from the material) when it is subjected to the reducing atmosphere. The overall reaction can be written as

For both undoped and Nb-doped samples, though the heating and cooling rates were identical, the total mass gain is higher than the total mass loss. This is consistent with the fact that the reduction process is slower than the re-oxidation one, as discussed in Sec. 3.5, and is therefore not complete during the TG analysis. Moreover, the total mass loss in the case of the Nb-doped sample is slightly larger than that for the undoped one. This could be due to the fact that the concentration of preexisting η– is higher in the former (see Sec. 3.3). Overall, the mass losses and gains are quite similar between undoped and Nbdoped samples. If we assume that the re-oxidation of the fully reduced samples is complete, the Ti3+/Ti4+ ratio can be estimated from the mass gain which is about 0.2 % for both samples. This is equivalent to about 0.1 oxygen and corresponds to the re-oxidation of about 6.4 % of Ti atoms from Ti3+ into Ti4+. The chemical formula of the fully reduced undoped material can, therefore, nominally be described as Ba3Ti2.84+Ti0.23+O5.9(BO3)2; here, again, we explicitly take into account only the defects that occur during the reduction process.

x x H 2  Ba 3Ti34 x Ti3x O6  x/2 (BO3 ) 2  H 2O (2) 2 2 Here, the positively charged oxygen vacancies VO2+ (i.e., the removal of O2–) are charge-compensated by negatively charged (electron) polarons η– (i.e., Ti3+ at the Ti4+ site). The removed O2– species then reacts with hydrogen to form water that is subsequently released. We note that, for simplicity, the defects that preexist in the material before the reduction process (see Fig. 6a) are not explicitly described in the chemical formulae in Eq. (2); i.e., only Ba 3Ti3O6 (BO3 ) 2 

Figure 10. Thermogravimetric analyses of (a) unreduced Ba3Ti3O6(BO3)2 in Ar/H2 5% and fully reduced Ba3Ti3O6(BO3)2 in O2 and (b) unreduced Ba3Ti3– xMxO6(BO3)2 in Ar/H2 5% and fully reduced Ba3Ti3– xMxO6(BO3)2 in O2. The measurements were carried out with same heating (from RT to 900°C) and cooling (from 900°C to RT) 1 K min-1 rates.

TG analyses also show a mass gain during the heating step of the fully reduced samples in air; see the red curves in Fig. 10. From RT to about 200°C, the sample masses stay almost the same; a rapid mass increase occurs between 200 and 300°C and followed by a smoother mass increase. The overall process can formally be described by the following re-oxidation reaction: Ba 3Ti34 xTi3x O6  x/2 (BO3 )2 

x O2  Ba 3Ti3O6 (BO3 )2 4

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In this case, oxygen from the atmosphere is incorporated into the fully reduced material and the oxygen vacancies are refilled. Charge neutrality is maintained through the conversion of the Ti3+ ions in the reduced samples back into Ti4+, i.e., the annihilation of both VO2+ and η–.

3.8. Determination of the Ti3+/Ti4+ Ratio and the Fermi-Level Position. XPS was employed to examine the oxidation state of titanium species and the electronic structure of the unreduced and fully reduced Ba3Ti3O6(BO3)2 and Ba3Ti3–xMxO6(BO3)2 (x = 0.06; M = Nb5+). Figure 11 shows the XPS spectra of Ti and Nb core levels. From the fitting of the peak area of the Ti core level by considering contributions from both Ti3+ and Ti4+, the Ti3+ content in the fully reduced samples was quantified. For the unreduced samples, the fit of the Ti4+/Ti3+ ratio gives negligible Ti3+ amount (i.e., the value is within the error bar of the fit), which indicates that only Ti4+ is present in the unreduced samples. Regarding the fully reduced samples, the Ti3+ content is about 4.0 % in the case of Ba3Ti3O6(BO3)2 and 5.3 % in the case of Ba3Ti3– xNbxO6(BO3)2 (x = 0.06). These results are in agreement with those determined by TG analysis (see Sec. 3.7). The XPS study shows that the Ti3+ content is slightly higher in the case of the Nb-doped sample than in the undoped one, which is again consistent with our analysis in Sec. 3.3. Nb5+ species was also detected by XPS for Nb-doped samples. The binding energy (BE) value of Nb5+ is the


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ACS Applied Energy Materials Figures 12a and 12b show the valence-band XPS of the abovementioned samples. The position of the VBM was obtained directly from the electron emission spectrum; the CBM position can then be deduced using the optical bandgap values reported in Sec. 3.1; the results are summarized in Fig. 12c. We find that the Fermi level of unreduced Ba3Ti3O6(BO3)2 is 2.22 eV above the VBM, in excellent agreement with the Fermi level μeint = 2.15 eV obtained in our calculations (see Ref. 14 and also Fig. 6a). As regards the fully reduced samples, the Fermi level is shifted toward the CBM which is, again, consistent with the results reported in Fig. 6b in which μeint,red is closer to the CBM compared to μeint in the undoped one. Regarding the effects of doping, the Fermi level of the Nb-doped samples is slightly shifted compared to that of the undoped ones; see Fig. 12c. This is consistent with the picture for NbTi presented in Fig. 6c in which the impurity, when incorporated into the material with a concentration higher than that of VO2+, can shift the Fermi level toward the CBM. 4. CONCLUSIONS

Figure 11. XPS spectra of unreduced and fully reduced (a) Ba3Ti3O6(BO3)2 showing Ti 2p core levels and (b) Ba3Ti3– 4+ has xNbxO6(BO3)2 showing Ti 2p and Nb 3d core levels. Ti a binding energy (BE) of 458.2 eV for the 2p3/2 transition in all samples, except in unreduced Ba3Ti3O6(BO3)2 where it has a BE of 458.3 eV; Ti3+ has a BE of 456.7 eV for the 2p3/2 transition in the fully reduced samples; and Nb5+ has a BE of 206.8 eV for the 3d5/2 transition in the Nb-doped samples. The BE values are in good agreement with those reported in the literature.27 Experimental data are represented by points and model fitting by solid lines.

same as that reported in Ti0.96Nb0.04O2 and Ti0.93Nb0.07O2.27 The XPS data does not indicate the presence of any Nb species with lower oxidation states, thus confirming the computational results in Sec. 3.3.

A detailed study of the stability and transport properties of undoped and doped Ba3Ti3O6(BO3)2 in Ar/H2 5% was presented. The material undergoes a change in the electrical transport mechanism, from being purely ionic to mainly electronic, when subjected to the extreme reducing atmosphere. The ionic-to-electronic transition is caused by a change in the defect landscape of the material according to which the highly mobile electron polarons (i.e., Ti3+, which is reduced from Ti4+) replace the low mobile oxygen defects as the predominant current-carrying species. The material is stable in the reducing atmosphere and the Ti4+/Ti3+ reduction and re-oxidation reactions are reversible. More importantly, by using Ba3Ti3O6(BO3)2 as an example, we have thus developed a novel theoretical framework and experimental procedures that can also be applied to understanding other SOFC materials. On the applications side, the stability and reversibility of

Figure 12. Valence-band XPS spectra of unreduced and fully reduced (a) Ba 3Ti3O6(BO3)2 and (b) Ba3Ti3–xNbxO6(BO3)2. (c) Schematic view of the positions of the VBM (lower part) and CBM (upper part) according to the XPS and optical diffuse reflectance data.


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the reduction/re-oxidation reactions are of great interest for SOFCs. If Ba3Ti3O6(BO3)2-based materials are to be used at the anode side, because of the materials' low sintering temperature (950°C), the anode can be integrated into the fuel-cell architecture at the end of the cell manufacturing process; i.e., the materials would be suitable for cathode-supported SOFCs. Besides, the conductivity may be further improved through substitution of Ti with another transition metal. Taking into account this potential improvement, the materials may be used in a barrier layer between the anode and the electrolyte to prevent chemical reactivity and cation diffusion. They may also be part of a composite at the anode side with a metallic element such as copper (cermet) or with another anode material (cercer). This combination may improve the stability and/or the catalytic activity towards hydrogen oxidation and reforming of hydrocarbon fuels.

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AUTHOR INFORMATION Corresponding Authors

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* [email protected] (K.H.), * [email protected] (E.Q.)

Notes The authors declare no competing financial interest.

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ACKNOWLEDGMENT The authors thank Etienne Janod and Benoit Corraze for helping with the Seebeck measurements. J.-M. D.'s Ph.D. thesis was financially supported by Ministère de l'Education Nationale, de l’enseignement supérieur et de la Recherche, France. The calculations were carried out using computing resources at the Center for Computationally Assisted Science and Technology at North Dakota State University.

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Ionic to Electronic Transport in Ba3Ti3O6(BO3)2 under Reducing

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